Method of pre-aging nitihf shape memory alloys and parts therefrom with uniform microstructures and superior properties

ABSTRACT

A method to produce a high strength NiTiHf alloy, a NiTiHfZr alloy or a NiTiZr alloy are disclosed. The alloys comprise less than about 10 atomic percent of Hf, Hf+Zr, or Zr, respectively. The alloys, devices containing the alloys and methods of producing the devices are also disclosed.

CROSS REFERENCE TO RELATED APPLICATIONS

This application claims priority and the benefit under 35 U. S.C. §119(e) to U.S. Provisional Patent Application Ser. No. 62/609,848 filedDec. 22, 2017, which is incorporated herein in its entirety byreference.

GOVERNMENT LICENSE RIGHTS

This invention was made with government support under grant numberDE-SP0022534 awarded by the Department of Energy. The Government hascertain rights in the invention.

FIELD OF THE INVENTION

The present invention relates to a two-step aging treatment comprised ofa low-temperature pre-aging heat treatment step, followed by a typicalhigh-temperature aging treatment, the latter similar to that establishedfor higher Hf content alloys. The alloys produced by this treatment andproducts containing these alloys and methods of producing the productscontaining the alloys are also aspects of the invention.

BACKGROUND

NiTi-based shape memory alloys (SMAs) have been successfully used inmany different fields of engineering because of their functionalproperties that include shape memory effect (SME) and superelasticity(SE). These properties are due to the occurrence of reversible,thermoelastic martensitic transformations. Because of inherentlimitations to binary NiTi alloys, NiTiHf (i.e. nickel (Ni), titanium(Ti), and Hafnium (Hf)) alloys have primarily been proposed for use inhigh-temperature aerospace and automotive actuation applications whereNiTi cannot perform, namely ambient environments greater than about 100°C., as Hf additions above about 10 atomic percent begin to increase thetransformation temperatures. Hence, most research to date has focused onthe shape memory and superelastic behavior of NiTiHf alloys with a highHf content (generally on the order of 15-25 atomic percent.) For thesehigh Hf containing materials, strength and shape memory behaviors arestrongly enhanced by the formation of a precipitate reinforcedmicrostructure, with the main strengthening precipitate commonly knownas the H-phase. These precipitates are formed through a basic aging heattreatment, which is an effective way to increase the matrix strength inthese alloys by forming fine precipitates that act as pinning sitesagainst the movement of dislocations while still allowing themartensitic transformation to occur nearly unimpeded. The currentpractice for these high Hf-containing alloys is to age the solutiontreated material at 500 to 600° C. for between 1 and 4 hours, with themost common ageing treatment being 550° C. for 3 hours. This typicalheat treatment is used to achieve high recoverable strain, highstrength, excellent superelastic behavior, and microstructural anddimensional stability at temperatures greater than 100° C. in Ni-rich(>50 atomic percent Ni) NiTiHf alloys containing ≥15 atomic percent Hf.However, use of a single aging treatment, similar to that used for thehigh Hf alloys, has not produced the same uniform microstructures orresulted in sufficient mechanical and functional properties in Ni-richNiTiHf alloys containing less than 10 atomic percent Hf.

SUMMARY

The present invention relates to a two-step aging treatment comprised ofa low-temperature pre-aging heat treatment step, followed by a typicalhigh-temperature aging treatment.

An aspect of the invention is a method to pre-age or heat treat low Hf(i.e. less than or equal to about 10 at. %) NiTiHf alloys, withcompositions in the range of about 50 to about 53 atomic percent Ni,about 1 to about 10 atomic percent Hf, less than about 5 total atomicpercent of incidental materials, which can include O, C, Fe, Mn, Cr, V,Nb, Mo, Ta, or W, or combinations thereof, and the balance of the alloybeing Ti. The method comprises treating the NiTiHf alloys at atemperature between about 150° C. and about 400° C., for greater thanabout 10 minutes, in some embodiments between about 5 min and about 10days, in some embodiments between about 10 minutes and 24 hours. Abovetemperatures of about 400° C., the method will produce primaryprecipitation, which is preferably avoided during this step. Thispre-aging step results in homogeneous clusters of Ni and/or Hf atomsthat act as nucleation sites for the precipitation of the H-phase,during subsequent aging at higher temperatures, resulting in a highstrength NiTiHf alloy with a homogenous distribution of precipitates.Without the pre-aging treatment, the nucleation of H-phase tends to beheterogeneous, nucleating on such defects as grain boundaries,dislocations, and other microstructural features resulting in aninhomogeneous distribution of large precipitates (e.g. where the spacingbetween precipitates is more than about 1.5 times the size of theprecipitates), which can be dependent upon the number of defects in theoriginal material and can vary with processing. This inhomogeneousdistribution of coarse precipitates results in poor mechanical andfunctional properties.

An aspect of the invention is a method to pre-age a NiTiHf alloy. Themethod comprises annealing the NiTiHf alloy at a temperature betweenabout 700° C. and about 1150° C., followed by a pre-aging step at atemperature between about 100° C. and about 400° C., for between about10 min and about 24 hours to produce a pre-aged NiTiHf alloy, andfinally heat treating the pre-aged NiTiHf alloy at a temperature ofbetween about 400° C. and about 600° C. to produce a homogenouslydistributed and high density of nano-sized H-phase precipitatesthroughout the matrix of a NiTiHf alloy. In some embodiments, theannealing can be solution annealing. The resulting alloy has improvedmechanical stability and strength and at least equal to or greater thanabout 4% recoverable compression strain, in some embodiments betweenabout 2% and 8% recoverable compression strain, without permanentdeformation. The absence of a pre-aging step generally results inheterogeneous formations of larger H-phase precipitates, primarilyclustered along grain boundaries, and correlated with poorer mechanicalbehavior.

An aspect of the invention is a method to produce a high-strengthNiTi(Hf/Zr) alloy with homogeneous distribution of H-phase precipitates.The recoverable compression strain is between about 2 and about 8%. Themethod includes processing a NiTi(Hf/Zr) alloy by at least one method ofcasting, additive manufacturing, drawing, forging, extrusion, powdermetallurgy, or combinations thereof. The composition of the NiTi(Hf/Zr)alloy includes between 50 and 53 atomic percent of Ni, between about 1and 10 atomic percent of Hf, Zr, or combinations thereof, less thanabout 5 atomic percent total% of incidental materials, and the balanceof the composition being Ti. The NiTi(Hf/Zr) workpiece is annealed toproduce an annealed NiTi workpiece, which is then pre-aged at atemperature between about 100° C. and about 400° C. for a pre-aging timeperiod between about 1 and 24 hours to produce the high-strength NiTialloy.

An aspect of the invention is a NiTi alloy. The recoverable yield of thealloy is between about 2 and about 8%, and the compressive yieldstrength of the alloy is greater than about 1.5 GPa.

An aspect of the invention is a biomedical implant. The implant includesa NiTi alloy, wherein a composition of the NiTi alloy isNi_(50.3-53.)Ti_(49.7-x)Hf_(x)Zr_(y) wherein x is between 6 and 9, 3 and6, or 1 and 3, wherein x+y is between 1 and 10, and wherein y is between0 and 10.

BRIEF SUMMARY OF THE DRAWINGS

The patent or application file contains at least one drawing executed incolor. Copies of this patent or patent application publication withcolor drawing(s) will be provided by the Office upon request and paymentof the necessary fee.

FIG. 1A illustrates the mechanical behavior of Ni_(50.3)Ti_(42.7)Hf₆(at. %) alloys in compression with pre-aging treatment at 300° C. for 12hours followed by normal aging at 550° C. for 3.5 hours for all samples.The solid arrows show the amount of strain recovered upon heating thedeformed samples to 150° C.;

FIG. 1B illustrates the mechanical behavior of Ni_(50.3)Ti_(41.7)Hf₈(at. %) alloys in compression with pre-aging treatment at 300° C. for 12hours followed by normal aging at 550° C. for 3.5 hours for all samples.The solid arrows show the amount of strain recovered upon heating thedeformed samples to 150° C.;

FIG. 1C illustrates the mechanical behavior ofNi_(50.3)Ti_(41.2)Hf_(8.5) (at. %) alloys in compression with pre-agingtreatment at 300° C. for 12 hours followed by normal aging at 550° C.for 3.5 hours for all samples. The solid arrows show the amount ofstrain recovered upon heating the deformed samples to 150° C.;

FIG. 1D illustrates the mechanical behavior of Ni_(50.3)Ti_(40.7)Hf₉(at. %) alloys in compression with pre-aging treatment at 300° C. for 12hours followed by normal aging at 550° C. for 3.5 hours for all samples.The solid arrows show the amount of strain recovered upon heating thedeformed samples to 150° C.;

FIG. 1E illustrates the mechanical behavior of Ni_(50.3)Ti_(42.7)Hf₆(at. %) alloys in compression without pre-aging treatment at 300° C. for12 hours, but including normal aging at 550° C. for 3.5 hours for allsamples. The solid arrows show the amount of strain recovered uponheating the deformed samples to 150° C.;

FIG. 1F illustrates the mechanical behavior of Ni_(50.3)Ti_(41.7)Hf₈(at. %) alloys in compression without pre-aging treatment at 300° C. for12 hours, but including normal aging at 550° C. for 3.5 hours for allsamples. The solid arrows show the amount of strain recovered uponheating the deformed samples to 150° C.;

FIG. 1G illustrates the mechanical behavior ofNi_(50.3)Ti_(41.2)Hf_(8.5) (at. %) alloys in compression withoutpre-aging treatment at 300° C. for 12 hours, but including normal agingat 550° C. for 3.5 hours for all samples. The solid arrows show theamount of strain recovered upon heating the deformed samples to 150° C.;

FIG. 1H illustrates the mechanical behavior of Ni_(50.3)Ti_(40.7)Hf₉(at. %) alloys in compression without pre-aging treatment at 300° C. for12 hours, but including normal aging at 550° C. for 3.5 hours for allsamples. The solid arrows show the amount of strain recovered uponheating the deformed samples to 150° C.;

FIG. 2A illustrates a conventional bright field image ofNi_(50.3)Ti_(41.2)Hf_(8.5) (at. %) after pre-aging at 300° C. for 12hours. Corresponding selected area diffraction pattern along [111] zoneaxis is illustrated in the lower left inset;

FIG. 2B illustrates a high resolution TEM micrograph along [111] zoneaxis of Ni_(50.3)Ti_(41.2)Hf_(8.5) (at. %) alloy pre-aged at 300° C. for12 hours and subsequently aged at 550° C. for 3.5 hours. Some of theH-phase precipitates are indicated by arrows marked with the letter “P”.The corresponding Fast Fourier transform (FFT) pattern is illustrated inthe lower right inset;

FIG. 3A illustrates conventional Bright Field (BF) images ofNi_(50.3)Ti_(42.7)Hf₆ (at. %) at 50 nm after aging at 550° C. for 3.5hours without the pre-aging treatment, with corresponding Selected AreaElectron Diffraction (SAED) patterns along [111] zone axis for eachregion presented in the inset;

FIG. 3B illustrates conventional BF images of Ni_(50.3)Ti_(42.7)Hf₆ (at.%) at 200 nm after aging at 550° C. for 3.5 hours without the pre-agingtreatment, with corresponding SAED patterns along [111] zone axis foreach region presented in the inset;

FIG. 3C illustrates conventional BF images of Ni_(50.3)Ti_(41.2)Hf_(8.5)(at. %) at 100 nm after aging at 550° C. for 3.5 hours without thepre-aging treatment, with corresponding SAED patterns along [111] zoneaxis for each region presented in the inset; and

FIG. 3D illustrates conventional BF images of Ni_(50.3)Ti_(40.7)Hf₉ (at.%) at 50 nm after aging at 550° C. for 3.5 hours without the pre-agingtreatment, with corresponding SAED patterns along [111] zone axis foreach region presented in the inset.

DETAILED DESCRIPTION

The present invention generally relates to a two-step aging treatment ofmetal alloys, and the material formed with this treatment. Zirconiumbehaves similarly as Hf in Ni-rich NiTi-based alloys forming the sameH-phase precipitates and consequently ternary and quaternary alloys withcompositions in the range of between about 50 and about 53 atomicpercent Ni, between about 1 and about 10 atomic percent Hf+Zr (wherezirconium (Zr) can range from about 0 to about 10 atomic percent), andthe balance being Ti and any incidental material (i.e. O, C, Fe, Mn, Cr,V, Nb, Mo, Ta, or W, and combinations thereof), would behave in asimilar manner as that described above, thus also benefitting from thepre-aging treatment disclosed.

The benefit of the pre-aging treatment is clearly demonstrated by thedata in Table 1. Table 1 includes mechanical data for various Ni-rich,NiTiHf alloys with and without the pre-aging treatment. The pre-agedsamples exhibited greater recoverable (and superelastic) strains whenloaded to about 2 GPa in compression and also did not undergo anypermanent (plastic) strain when loaded to that level, indicating thatthe yield strength of the material was greater than about 2 GPa. Table 1illustrates the benefits of the pre-aging treatment to enhance thestrength, and improve recoverable and superelastic strains of thevarious alloys compared to the same material without the pre-agingtreatment. Each value in Table 1 is approximate.

TABLE 1 Recoverable Permanent strain at strain at 2 GPa Composition (at.%) Heat treatment 2 GPa (%) (%) Ni_(51.5)Ti_(42.5)Hf₆ With pre-aging 4.30 Without pre aging 1.8 0 Ni₅₁Ti₄₁Hf₈ With pre-aging 4.75 0 Withoutpre-aging 2.5 0 Ni_(50.3)Ti_(43.7)Hf₆ With pre-aging 3.8 0 Without preaging 2.4 5 Ni_(50.3)Ti_(41.7)Hf₈ With pre-aging 4 0 Without pre aging3.4 3 Ni_(50.3)Ti_(41.2)Hf_(8.5) With pre-aging 4 0 Without pre aging 42

One aspect of the present invention is a method to produce ahigh-strength NiTiHf alloy with homogeneous distribution of H-phaseprecipitates, where the alloy comprises greater than or equal to about4% recoverable compression strain. The method includes processing analloy by at least one process of casting, additive manufacturing,drawing, forging, extrusion, powder metallurgy, or other typicalmetallurgical process, or combinations thereof. The composition of thealloy includes between about 50 and about 53 atomic percent Ni, betweenabout 1 and about 10 atomic percent of Hf, less than about 1 atomicpercent of incidental materials (e.g. O, C, Fe, Mn, Cr, V, Nb, Mo, Ta,or W, or combinations thereof), and the balance being Ti. Following theprocessing step, the alloy is solution treated at 700° C. and about1150° C. followed by pre-aging at a temperature between about 100° C.and about 400° C. for at least about 10 min.

In some embodiments, the temperature of the pre-aging step can bebetween about 300° C. and about 400° C., or about 200° C., about 250°C., about 300° C., about 350° C., about 400° C., about 450° C. The alloycan be maintained at the temperature for between about 10 min and 24hours, in some embodiments between about 3 hours and about 18 hours,between about 5 hours and about 12 hours, or between about 7 hours andabout 10 hours. In some embodiments, the alloy can be maintained at thepre-aging temperature for about 10, min, 30 min, one hour, about twohours, about three hours, about four hours, about five hours, about sixhours, about seven hours, about eight hours, about nine hours, about 10hours, about 11 hours, about 12 hours, about 13 hours, about 14 hours,about 15 hours, about 16 hours, about 17 hours, about 18 hours, about 19hours, about 20 hours, about 21 hours, about 22 hours, about 23 hours,or about 24 hours.

The NiTiHf alloy can comprise several compositions. In some embodiments,the alloy comprises Ni_(50.3-53)Ti_(49.7-x)Hf_(x) wherein x is between 6and 9, 3 and 6, or 1 and 3. Incidental materials can be included in theNiTiHf alloy, which can include up to about 1 atomic percent of anytransition metals (individually) or up to about 1 atomic % of a nonmetalsuch as oxygen or carbon, and combinations thereof. The maximum amountof the incidental materials can be about up to about 5 atomicpercentage. In some embodiments, the amount of incidental materials canbe between about 1 atomic percentage and about 5 atomic percentage. Insome embodiments, the Hf can be a combinations of Hf and Zr, wherein thesum of Hf+Zr is between about 1 and about 10 atomic percentage andwherein the contribution of Zr can be between about 0.1 and about 10atomic percentage (with 10 atomic percentage resulting inNi_(50.3-53)Ti_(49.7-x)Zrx).

The method can further include an annealing step prior to the pre-agingstep. The annealing temperature is between about 700° C. and about 1150°C. In some embodiments, the annealing temperature can be between about750° C. and about 1100° C., about 800° C. and about 1050° C., about 850°C. and about 1000° C., or about 900° C. and about 950° C. In someembodiments, the annealing temperature can be about 700° C., about 750°C., about 800° C., about 850° C., about 900° C., about 950° C., about1000° C., about 1050° C., about 1100° C., or about 1100° C. The alloycan be maintained at the annealing temperature for between about 30minutes and about 144 hours. In some embodiments, the annealing durationcan be about 72 hours for cast material. Following annealing, the alloycan be quenched at a temperature between about −100° C. and about 200°C.

In some embodiments, the method can further include a heat treatmentfollowing the pre-aging step. The pre-aged alloy can be heat treated ata temperature between about 400° C. and about 600° C. for between about10 minutes and about 24 hours. In some embodiments, the heat treatmenttemperature can be between about 500° C. and about 550° C.

An aspect of the invention is a method to produce a NiTiHf alloy. Themethod includes processing an alloy by at least one method of casting,additive manufacturing, drawing, forging, extrusion, powder metallurgy,or other typical metallurgical process, or combinations thereof. Thecomposition of the NiTiHf alloy is between about 50 and about 53 atomicpercent or Ni, between about 1 and about 10 atomic percent of Hf, nogreater than about 2% atomic percent (total) of incidental materials,and the balance being Ti. The NiTiHf alloy is solution annealed at atemperature between about 700° C. and about 1150° C. to produce asolution annealed NiTiHf alloy. The solution annealed NiTiHf alloy ispre-aged at a pre-aging temperature, which is between about 100° C. andabout 400° C. for a time period between about 1 hour and about 24 hoursto produce a pre-aged NiTiHf workpiece. The pre-aged NiTiHf workpiece isheat treated at a heat treatment temperature between about 400° C. andabout 600° C. to produce a high density of homogeneously distributedH-phase precipitates in the NiTiHf part or form.

In some embodiments, the temperature of the pre-aging step can bebetween about 300° C. and about 600° C., between about 400° C. and about550° C., or about 200° C., about 250° C., about 300° C., about 350° C.,about 400° C., about 450° C., about 500° C., about 600° C., about 650°C., or about 700° C. The alloy can be maintained at the temperature forbetween about one hour and 24 hours, in some embodiments between about 3hours and about 18 hours, between about 5 hours and about 12 hours, orbetween about 7 hours and about 10 hours. In some embodiments, the alloycan be maintained at the pre-aging temperature for about one hour, abouttwo hours, about three hours, about four hours, about five hours, aboutsix hours, about seven hours, about eight hours, about nine hours, about10 hours, about 11 hours, about 12 hours, about 13 hours, about 14hours, about 15 hours, about 16 hours, about 17 hours, about 18 hours,about 19 hours, about 20 hours, about 21 hours, about 22 hours, about 23hours, or about 24 hours.

The annealing temperature is generally between about 700° C. and about1150° C. In some embodiments, the annealing temperature can be betweenabout 750° C. and about 1100° C., about 800° C. and about 1050° C.,about 850° C. and about 1000° C., or about 900° C. and about 950° C. Insome embodiments, the annealing temperature can be about 700° C., about750° C., about 800° C., about 850° C., about 900° C., about 950° C.,about 1000° C., about 1050° C., about 1100° C., or about 1100° C. Thealloy can be maintained at the annealing temperature for between about30 minutes and about 144 hours, in some embodiments about 72 hours.Following annealing, the alloy can be quenched at a temperature betweenabout −100° C. and about 200° C.

The pre-aged alloy can be heat treated at a temperature between about400° C. and about 600° C. for between about 30 minutes and about 24hours. In some embodiments, the heat treatment temperature can bebetween about 500° C. and about 550° C.

The NiTiHf alloy can comprise several compositions. In some embodiments,the alloy comprises Ni_(50.3-53)Ti_(49.7-x)Hf_(x) wherein x is between 6and 9, 3 and 6, or 1 and 3. Incidental materials can be included in theNiTiHf alloy, which can include up to about 1 atomic percent of anytransition metals (individually) or up to about 1 atomic % of a nonmetalsuch as oxygen or carbon and combinations thereof. The maximum amount ofthe incidental materials can be about 5 atomic percentage. In someembodiments, the amount of incidental materials can be between about 1atomic percentage and about 5 atomic percentage. In some embodiments,the composition can be Ni_(50.3)Ti_(49.7-x)Hf_(x), where x can bebetween 6 and 9, in some embodiments x can be 6, 8, 8.5 or 9.

An aspect of the invention is a method to produce a NiTiHfZr alloy. Themethod includes processing an alloy workpiece by a method from the groupconsisting of casting, additive manufacturing, drawing, forging,extrusion, powder metallurgy, or other typical metallurgical process orcombinations thereof. The NiTiHfZr alloy comprises between about 50 andabout 53 atomic percent of Ni, between about 1 and 10 atomic percentHf+Zr (where Zr can range from about 0 to about 10 atomic percent), lessthan about 2 atomic percentage of incidental materials, and the balancebeing Ti. The NiTiHfZr workpiece is solution annealed at a temperatureof between about 700° C. and about 1150° C. to produce a solutionannealed NiTiHfZr workpiece. The annealed NiTiHfZr alloy is pre-aged ata pre-aging temperature of between about 100° C. and about 400° C. forbetween about 1 hour and about 24 hours to produce a pre-aged NiTiHfZralloy. The pre-aged NiTiHf Zr alloy is heat treated at a temperaturebetween about 400° C. and about 600° C. to produce a high density ofhomogeneously distributed H-phase precipitates in the NiTiHfZr alloy.

In some embodiments, the temperature of the pre-aging step can bebetween about 100° C. and about 300° C., between about 300° C. and about350° C., or about 200° C., about 250° C., about 300° C., about 350° C.,or about 400° C. The alloy can be maintained at the temperature forbetween about one hour and 24 hours, in some embodiments between about 3hours and about 18 hours, between about 5 hours and about 12 hours, orbetween about 7 hours and about 10 hours. In some embodiments, the alloycan be maintained at the pre-aging temperature for about one hour, abouttwo hours, about three hours, about four hours, about five hours, aboutsix hours, about seven hours, about eight hours, about nine hours, about10 hours, about 11 hours, about 12 hours, about 13 hours, about 14hours, about 15 hours, about 16 hours, about 17 hours, about 18 hours,about 19 hours, about 20 hours, about 21 hours, about 22 hours, about 23hours, or about 24 hours.

The annealing temperature is between about 700° C. and about 1150° C. Insome embodiments, the annealing temperature can be between about 750° C.and about 1100° C., about 800° C. and about 1050° C., about 850° C. andabout 1000° C., or about 900° C. and about 950° C. In some embodiments,the annealing temperature can be about 700° C., about 750° C., about800° C., about 850° C., about 900° C., about 950° C., about 1000° C.,about 1050° C., about 1100° C., or about 1100° C. The alloy can bemaintained at the annealing temperature for between about 30 minutes andabout 144 hours, in some embodiments about 72 hours. Followingannealing, the alloy can be quenched at a temperature between about−100° C. and about 200° C.

The pre-aged alloy can be heat treated at a temperature between about400° C. and about 600° C. for between about 1 minutes and about 24hours.In some embodiments, the heat treatment temperature can be between about500° C. and about 550° C.

An aspect of the invention is a NiTiHf alloy, NiTiZr alloy or a NiTiHfZralloy, wherein the alloy comprises a homogenously distributed H-phase inthe precipitate.

In some embodiments, the compression yield strength of the NiTiHf alloycan be greater than about 1.5 GPa. In some embodiments, the compressiveyield strength can be between about and about 8% recoverable compressionstrain. In some embodiments, the compression yield strength can begreater than about 2 GPa when tested at about room temperature. In someembodiments, the compressive strength can be up to about 4 GPa, in someembodiments between about 1.5 GPa and 4 GPa. The recoverable compressionstrain can be between about 2% and about 8%, in some embodiments about4%. The tensile strength can be between about 60% and 70% of thecompressive strength.

An aspect of the invention is a method for forming a biomedical implant.The method includes providing a NiTiHfZr alloy, a NiTiZr alloy or aNiTiHf alloy; and producing the biomedical implant by additivemanufacturing.

An aspect of the invention is a biomedical implant comprising a NiTiHfZralloy, NiTiZr alloy or NiTiHf alloy.

EXAMPLES Example 1

NiTiHf alloys with target compositions of Ni_(50.3)Ti_(50-x)Hf_(x), withx=6, 8, 8.5, and 9 at. % were made by induction-melting high-purityelemental constituents using a graphite crucible and casting into acopper mold. The ingots were homogenized in vacuum at 1050° C. for 72hours and then extruded at about 900° C. at a 7:1 area reduction ratio.The extruded rods were sectioned into samples that were initiallysolution-annealed at 1050° C. for 30 min, water quenched, and thenpre-aged at 300° C. for 12 h and air-cooled, and finally aged a secondtime at 550° C. for 3.5 hours and air-cooled. To isolate the effect ofpre-aging on the functional properties of NiTiHf alloys, other testsamples were directly aged at 550° C. for 3.5 hours after thesolution-anneal treatment (without pre-aging at 300° C. for 12 hours).

Mechanical compression tests were performed on an MTS servo-hydraulicload-frame equipped with an MTS 661.20 load cell. Compression sampleswere cylindrical with a diameter of 5 mm and a length of 10 mm. Fivecompression cycles were applied to the samples using a maximum load of40 kN and a minimum load of 250 N, corresponding to 2 GPa and 13 MPaengineering stress limits. A cross-head speed of 0.1 mm/min was used,corresponding to an approximate strain rate of 10-4 s⁻¹. The surfaces ofthe samples were speckled using an airbrush to deposit sequential layersof alumina powders (≤10 μm) and Brother TB450 carbon black toner powder.Digital images were acquired during loading using a Basler ACA2500camera, and the Ncorr digital image correlation (DIC) software was usedto analyze the displacements of the particles. From these displacementfields, the software calculated the surface strains during deformation.Before each test, eight images of the undeformed sample were acquiredand analyzed to establish the strain noise for each pattern, which fellbetween 10⁻⁴ to 10⁻⁵ for the data in FIG. 1. Because these figures arecompared to each other, they are provided as a single figure withsub-figures.

Conventional and high-resolution transmission electron microscopy (HRTEMor TEM) of aged NiTiHf samples was carried out using an FEI Talos TEM(FEG, 200 kV). The TEM samples were prepared by grinding slices to athickness of 90-100 μm; a mechanical punch was then used to create discswith a diameter of 3 mm. A Fischione automatic twin-jet electropolisher(model 120) at 13 V and an electrolyte of 30% HNO₃ in methanol (byvolume) at around −35° C. was then used to further thin the TEM foils.To measure the size of H-phase precipitates and interparticle distance(the distance of a single precipitate from its closest precipitate),several HRTEM images taken from various regions, were used. Thismeasurement was repeated for almost 100 precipitates on each sample andaverage precipitate size, average interparticle distance and theircorresponding standard error is reported.

FIG. 1 illustrates the compression responses of the NiTiHf alloys agedat 550° C. for 3.5 hours, with (FIGS. 1A-1D) and without (FIGS. 1E-1H)an initial pre-aging treatment of 300° C. for 12 hours. The compressiontests were performed at room temperature (about 23° C.), and all thesamples were initially austenite. The mechanical loading stress-inducedmartensite that did not recover upon unloading at room temperature.Heating the samples to about 150° C. and measuring the recovered strainsassessed how much of the deformation was due to stress inducedmartensite vs. plasticity, in a phenomenological sense. Completerecovery of the strain (almost 4%) was observed for all the NiTiHfsamples that were pre-aged at 300° C. for 12 hour (solid arrows in FIG.1A-1D). For the other samples with 6 to 8.5 at. % Hf (FIGS. 1E-1F), onlya portion of the strain was recovered upon heating, indicating astronger presence of plasticity in the mechanical responses sanspre-aging. For the Ni_(50.3)Ti_(40.7)Hf₉ sample (FIG. 1G), however, thestrain was fully recovered sans pre-aging, as well as in the pre-agedcase. While superelastic behavior was not directly observed, these datado suggest that the pre-aged samples and the un-pre-aged 9Hf samplewould likely exhibit stable superelastic responses at higher testtemperatures.

FIG. 2A illustrates a representative bright field (BF) TEM image of theNi_(50.3)Ti₄L₂Hf_(8.5) sample after pre-aging at 300° C. for 12 hours.Fully formed H-phase precipitates were not observed. The correspondingselected area electron diffraction pattern (SAED) of the B2 zone (upperright inset in FIG. 2A) confirms this observation, as the characteristicsuper-reflection of H-phase precipitates along <110> directions were notdetected; however, a clear structured pattern of diffuse intensity ispresent. This pattern of diffuse intensity was also observed for othersamples with different Hf content after pre-aging at 300° C. for 12hours. A HRTEM study of the sample further corroborated the lack offully formed H-phase precipitates after pre-aging at 300° C. for 12 hour(as demonstrated in the lower left inset in FIG. 2A).

These diffuse intensities with periodical character in reciprocal spaceindicate the existence of short-range order in the real-space lattice.Such diffuse intensities have been reported in binary Ni_(50.6)T_(149.4)after low-temperature aging and were attributed to the existence ofmicrodomains in the form of clusters of Ni atoms as precursors to fullformation of Ni₄Ti₃ nanoprecipitates. Analysis of H-phase precipitatesin NiTiHf alloys indicates that the Hf content of the H-phase is higherthan the Hf content in the matrix while the Ni content is also slightlyhigher in the precipitate compared with the matrix. It is thereforeexpected that Hf and/or Ni atom clusters form after low-temperatureaging (pre-aging treatment) as a precursor to H-phase precipitation uponfurther aging at higher temperatures, similar to the sequence ofmechanisms that give rise to Ni₄Ti₃ nano-precipitation in binary NiTialloys.

FIG. 2B is a HRTEM micrograph of the Ni_(50.3)Ti₄L₂Hf_(8.5) alloypre-aged at 300° C. for 12 hours and subsequently aged at 550° C. for3.5 h. The corresponding fast Fourier transform (FFT) is illustrated inthe upper left inset. The primary spots in the FFT result from the B2cubic austenite structure, and the super reflections at ⅓ positionsalong <110> B2, are reflections from uniquely oriented H-phaseprecipitates (depicted in FIG. 2B with the letter “P”). The H-phaseprecipitate dimensions for Ni50.3Ti49.7-xHfx alloys (where x is 6, 8,8.5, and 9 atomic percentage, respectively) with and without pre-agingat 300° C. for 12 hours followed by normal aging at 550° C. for 3.5hours are depicted in Table 2. For the latter cases (without pre-aging),characteristic precipitate morphologies are described for both regionsat grain boundaries (GB) as well as within the grain interior (IG). Allvalues in Table 2 are approximate.

TABLE 2 Without pre-aging With pre-aging Ni_(50.3)Ti_(42.7)Hf₆Ni_(50.3)Ti_(42.7)Hf₆ Ni_(50.3)Ti_(41.2)Hf_(8.5)Ni_(50.3)Ti_(41.2)Hf_(8.5) Ni_(50.3)Ti_(40.7)Hf₉ All alloys (GB) (IG)(GB) (IG) (GB & IG) Length 14 ± 1  60 ± 4 110 ± 10 43 ± 1 93 ± 2 23 ± 1(nm) Width 7 ± 1 15 ± 2 57 ± 5 12 ± 1 25 ± 1 10 ± 1 (nm) Inter- 8 ± 1  5± 1 95-700  5 ± 1 100 ± 4  12 ± 1 particle distance (nm)

For all the pre-aged samples, regardless of Hf content, the precipitateswere ellipsoidal in shape with average dimensions of 14±1 nm (length)and 7±1 nm (width); the interparticle distance was 8±1 nm. Obviousvariation in H-phase precipitate morphologies were not detected for anyof the pre-aged NiTiHf alloys. Therefore, the fully reversible 4%deformation depicted in FIGS. 1A-1D is aided by the uniform distributionof finely spaced H-phase precipitates, which strengthen the matrixagainst plastic deformation, while not impeding the formation ofmartensite.

The microstructure of the Ni_(50.3)Ti_(42.7)Hf₆ alloy that was aged at550° C. for 3.5 hours (without pre-aging) is illustrated in the BF-TEMmicrographs of FIGS. 3A and 3B taken from the vicinity of a grainboundary (GB) and the interior of the austenite grain, respectively. Itis apparent that aging of the Ni_(50.3)Ti_(42.7)Hf₆ alloy at 550° C. for3.5 hours induces heterogeneous nucleation of spindle-like H-phaseprecipitates along GBs (with an aspect ratio of 4) and a sparsedistribution of larger, longer, and widely spaced ellipsoidal H-phaseprecipitates (with an aspect ratio of 2) in the interior of the grains,though the major part of the grain shown in FIG. 3B is actually free ofprecipitates. GBs are well known locations for heterogeneous nucleationin solid state precipitation processes because they decrease theinterfacial energy between the precipitate and the parent phase, reducestress fields, and because of the chemical composition gradient thatexists around GBs. However, because there is competitive growth on theGBs, the size of the grain boundary precipitates is small compared withthe precipitates formed in the grain interiors.

With an increase in Hf content, the interparticle distance and size ofthe precipitates formed on the GB's and inside the grains decreases, asillustrated in FIG. 3C for the Ni_(50.3)Ti_(41.2)Hf_(8.5) sample (Table2). This suggests that there is transition from heterogeneousprecipitation in the 550° C. aged NiTiHf alloys (for low Hf content) tohomogeneous precipitation in alloys with higher Hf content for the sameNi:(Ti+Hf) ratio. The microstructure of Ni_(50.3)Ti_(40.7)Hf₉ afteraging at 550° C. for 3.5 hours (without pre-aging) is illustrated inFIG. 3D.

Clearly, the H-phase precipitates are homogenously nucleated anddistributed in grain interiors and regions near GBs; heterogeneous GBprecipitation is not observed. Consistent with the microstructuralobservations, the amount of permanent strain observed during compressiontesting decreased as the Hf content increased from 5.7% (FIG. 1E forNi_(50.3)Ti_(42.7)Hf₆) to 2% (FIG. 1G for Ni_(50.3)Ti_(41.2)Hf_(8.5)),and finally, full recovery without permanent deformation was observedfor the Ni_(50.3)Ti_(40.7)Hf₉ alloy, where precipitation was homogeneousand finely spaced.

Moreover, the H-phase precipitate dimensions of Ni_(50.3)Ti_(40.7)Hf₉alloy, with the application of pre-aging treatment (Table 2), areslightly smaller as is the interparticle distance compared to the samealloy only aged at 550° C. for 3.5 hours (without pre-aging). Therefore,the lower critical plateau stress of the Ni_(50.3)Ti_(40.7)Hf₉ samplewithout the pre-aging treatment (300 MPa in FIG. 1H) compared with thesame material with pre-aging treatment (520 MPa in FIG. 1D) could be dueto a larger interparticle distance where martensite can propagate moreeasily between particles. Other possibilities include a differentH-phase precipitate chemistry between the two conditions, resulting indifferent matrix chemistries, hence different transformationtemperatures, and/or the slightly larger coherent precipitates causelarger coherency strains, hence more stress and lower transformationstresses.

Finally, the stresses required to form martensite (the plateau stresses)for the samples that were pre-aged at 300° C. for 12 hours wererelatively consistent, falling between about 410 and 510 MPa (FIGS.1A-1D), yet in contrast, the martensite formation plateau stresses forthe samples without the pre-aging treatment (FIG. 1E-1H) continuouslydecreased from about 810 MPa to about 300 MPa as Hf content increasedfrom 6 to 9 at. %. In the former case, it is possible that the chemicaldiffusion mechanism that is promoted by pre-aging also promotes moreavailable hafnium to go into H-phase precipitates as the overallHf-content is increased for each different composition. In other words,the extra time for diffusion results in an increase in the Hf contentthat goes into the precipitates themselves, but a relatively uniform Hfand/or Ni content left in the matrix (or at least ratios that balanceeach other in terms of transformation temperature effects). Contrarily,the changes observed in the plateau stresses in the latter case (nopre-aging) are logically expected from the increase in overall Hfcontent of each alloy, given that for most of the compositions, H-phaseprecipitation is not uniform enough throughout the matrix to effect thetransformation temperatures greatly. Increasing the Hf content in NiTiHfalloys increases the Martensite start (Ms) temperature, thus through theClausius-Clapeyron relation, the martensite formation stresses woulddecrease given a constant test temperature. In the higher Hf contentswhere there is relatively uniform H-phase precipitation, the combinedmechanisms discussed in the previous paragraph likely all interact tostill cause a relative decrease in transformation temperatures relativeto the lower Hf alloys.

Pre-aging Ni_(50.3)Ti_(429.7-x)Hf_(x) alloys (x=6, 8, 8.5, 9 at. %) atabout 300° C. for about 12 hours after solution annealing at about 1050°C. leads to a higher density of the H-phase precipitates duringsubsequent aging at 550° C., than aging alone. This results in greaterstrength and subsequently improved mechanical and functional performancewith 4% recoverable compression strain. The results suggest that Hf orother atom clusters form during the pre-aging treatment resulting innucleation of a uniform distribution of H-phase when subsequently agedat 550° C. Aging NiTiHf alloys at 550° C. directly after the solutionannealing treatment (without pre-aging), leads to heterogeneousnucleation of H-phase precipitates on the grain boundaries in the lowerHf compositions (6 to 8 at. %). Most of the grain interior is free ofthe precipitates, which results in reduced strength and leads to poormechanical behavior. However, as Hf content is increased, homogenousprecipitation is achieved directly, without pre-aging. As a result,mechanical and functional behavior is improved compared to thecompositions with lower Hf content and no pre-aging.

Ranges, for example temperature ranges, atomic percentages, and others,have been discussed and used within the forgoing description. Oneskilled in the art would understand that any sub-range within the statedrange would be suitable, as would any number within the broad range,without deviating from the invention.

The foregoing description of the present invention related to NiTiHfalloys has been presented for purposes of illustration and description.Furthermore, the description is not intended to limit the invention tothe form disclosed herein. Consequently, variations and modificationscommensurate with the above teachings, and the skill or knowledge of therelevant art, are within the scope of the present invention. Theembodiment described hereinabove is further intended to explain the bestmode known for practicing the invention and to enable others skilled inthe art to utilize the invention in such, or other, embodiments and withvarious modifications required by the particular applications or uses ofthe present invention. It is intended that the appended claims beconstrued to include alternative embodiments to the extent permitted bythe prior art.

1. A method to produce a high-strength NiTi(Hf/Zr) alloy withhomogeneous distribution of H-phase precipitates, with between about 2and about 8% recoverable compression strain, comprising: processing aNiTi(Hf/Zr) alloy by at least one method of a casting method, anadditive manufacturing method, a drawing method, a forging method, anextrusion method, a powder metallurgy method, or combinations thereof,wherein the composition of the NiTi(Hf/Zr) alloy comprises between 50and 53 atomic percent of Ni, between about 1 and 10 atomic percent ofHf, Zr, or combinations thereof, less than about 5 atomic percent total% of incidental materials, and the balance of the composition being Ti;annealing the NiTi(Hf/Zr) workpiece to produce an annealed NiTiworkpiece; and pre-aging the annealed NiTi workpiece at a temperaturebetween about 100° C. and about 400° C. for a pre-aging time periodbetween about 1 and 24 hours to produce the high-strength NiTi alloy. 2.The method of claim 1, wherein the pre-aging temperature is about 300°C.
 3. The method of claim 2, wherein the pre-aging time period is about12 hours.
 4. The method of claim 1, further comprising quenching theannealed NiTi workpiece at a temperature between about −100° C. andabout 200° C. prior to pre-aging.
 5. The method of claim 1, wherein acompressive yield strength of the high-strength NiTi alloy is betweenabout 2 GPa and about 4 GPa.
 6. The method of claim 1, furthercomprising a heat treatment at a heat treatment temperature of betweenabout 400° C. and about 600° C. for between about 10 minutes and about24 hours.
 7. The method of claim 1, wherein the annealing is solutionannealing.
 8. The method of claim 7, wherein the solution annealingtemperature is between about 700° C. and about 1150° C.
 9. The method ofclaim 1, wherein a primary precipitate is not formed following thepreaging step.
 10. The method of claim 1, wherein the NiTi alloy is ofthe formula Ni50+x(Ti+Hf)49.9-1x wherein x is between 0.1 and 2.9. 11.The method of claim 10, wherein Hf in the NiTiHf alloy is between about0.1 to 10%.
 12. The method of claim 1, wherein the NiTi alloy is of theformula Ni50+x(Ti+Zr)49.9-1x wherein x is between 0.1 and 2.9.
 13. Themethod of claim 10, wherein Zr in the NiTi alloy is between about 0.1 to10%.
 14. The method of claim 10, wherein the NiTi alloy furthercomprises Zr, wherein the Zr in the NiTi alloy is between about 0 andabout 10 atomic percentage of a percentage of Hf in the NiTiHf alloy ofa composition NiTiHfZr.
 15. A NiTi alloy, wherein a recoverable yield ofthe alloy is between about 2 and about 8%, and wherein a compressiveyield strength of the alloy is greater than about 1.5 GPa.
 16. The NiTialloy of claim 15, wherein a compressive yield strength of thehigh-strength NiTi alloy is between about 2 GPa and about 4 GPa.
 17. TheNiTi alloy of claim 15, wherein a composition of the NiTi alloy isNi_(50.3-53.)Ti_(49.7-x)Hf_(x) wherein x is between 6 and 9, 3 and 6, or1 and
 3. 18. The NiTi alloy of claim 15, further comprising incidentalmaterials of up to about 1 atomic percent of any individual transitionmetals, up to about 1 atomic % of a nonmetal, and combinations thereof,wherein a total maximum amount of the incidental materials is about 5atomic percentage.
 19. A biomedical implant, comprising a NiTi alloy,wherein a composition of the NiTi alloy isNi_(50.3-53.)Ti_(49.7-x)Hf_(x) wherein x is between 6 and 9, 3 and 6, or1 and 3.